Advanced case carburizing secondary hardening steels

ABSTRACT

Steel alloys susceptible to case and core hardening comprising 0.05 to 0.24 weight percent carbon; 15 to 28 weight percent cobalt and 1.5 to 9.5 weight percent in nickel, small percentages of one or more additives: chromium, molybdenum, and vanadium; and the balance iron.

BACKGROUND OF THE INVENTION

[0001] This invention relates to a new class of steel alloys especiallyuseful for the manufacture of case hardened gears and other productsmade from case carburized steel alloys.

[0002] Currently, there are a number of high performance gear andbearing steels on the market. A number of these materials utilizeprimary carbides to achieve their high surface hardness and others usestage one or stage three tempered conditions with epsilon carbide orcementite strengthening. Primary carbides are formed when the carboncontent exceeds the solubility limit during the solution treatment andlarge alloy carbides precipitate. This is the case in particular forsecondary hardening steels using alloy carbide strengthening for greaterthermal stability to improve properties such as scoring resistance.However, research indicates that primary carbide formation can have adetrimental impact on both bending and contact fatigue resistance.Formation of primary carbides can also make process control difficultfor avoidance of undesirable carbide distributions such as networks. Inaddition, primary carbide formation in current gear and bearing steelcan lead to a reversal in the beneficial residual compressive stressesat the surface. This is due to a reversal of the spatial distribution ofthe martensite start temperature due to the consumption of austenitestabilizing elements by the primary carbides. Thus, there has developeda need for case hardenable steel alloys which do not rely upon primarycarbide formation, but provide secondary hardening behavior for superiorthermal stability. This invention provides a new class of steel meetingthis requirement, while exploiting more efficient secondary hardeningbehavior to allow higher surface hardness levels for even greaterimprovements in fatigue and wear resistance.

[0003] In applications of sliding wear the formation of primary carbidescan be beneficial; however, in current gear and bearing steels this canlead to a reversal in the beneficial residual compressive stresses atthe surface due to the consumption of elements promoting hardenabilityby the primary carbides.

[0004] Thus, there has developed a need for case hardenable steel alloyswhich do not rely upon primary carbide formation.

SUMMARY OF THE INVENTION

[0005] Briefly, the present invention comprises a class of casehardenable steel alloys with carbon content in the range of about 0.05weight percent to about 0.24 weight percent in combination with amixture of about 15 to 28 weight percent cobalt, 1.5 to 9.5 weightpercent nickel, 3.5 to 9.0 weight percent chromium, up to 3.5 weightpercent molybdenum, and up to 0.2 weight percent vanadium.

[0006] The microstructural features are a Ni—Co lath martensite matrixsteel strengthened by M₂C carbides typically containing Cr, Mo and V.Typical processing of this class of steels includes case carburizing,solution treatment, quenching, and tempering, although due to the highalloy content, quenching may not be required. Case carburizing producesa gradient in the volume fraction of the M₂C carbides and results in aconcomitant increase in hardness and promotes a surface residualcompressive stress. The efficiency of the M₂C strengthening responseallows this class of steels to achieve very high surface hardnesses withlimited soluble carbon content. Thus, this class of steels have theability to achieve very high surface hardnesses without the formation ofprimary carbides.

[0007] Typical advantages of this class of alloys include ultrahigh casehardness leading to superior wear and fatigue resistance, superior corestrength and toughness properties, optional air hardening resulting inless distortion, and higher thermal resistance.

[0008] This new class of secondary hardening gear and bearing steels arematrix steels utilizing an efficient M₂C precipitate strengtheningdispersion. Because of the efficiency of this strengthening dispersion,a superior combination of properties can be attained for a givenapplication. For example, in situations where the desired surfaceproperties are similar to current materials, the core strength andtoughness can be superior. In applications where superior surfaceproperties are desired, the disclosed steels can easily outperformtypical materials while maintaining normal core properties, and inapplications which require corrosion resistance, these new steels canprovide stainless properties with surface mechanical properties similarto typical non-stainless grades.

[0009] These and other objects, advantages and features of the inventionwill be set forth in the detailed description which follows.

BRIEF DESCRIPTION OF THE DRAWING

[0010] In the detailed description which follows, reference will be madeto the drawing comprised of the following figures:

[0011]FIG. 1 is a graph correlating hardness to precipitation drivingforce for experimental and predicted results;

[0012]FIG. 2 is a graph correlating precipitation half completion timeand half completion coarsening rate constant for experimental andpredicted results;

[0013]FIG. 3 is a graph which correlates calculated segregation freeenergy difference with experimental embrittlement potency;

[0014]FIG. 4 is a flow block diagram of the total system structure ofthe alloys of the invention;

[0015]FIG. 5 is a graph depicting the relationship between cobalt andnickel content for a 200° C. Ms temperature for the alloys of theinvention;

[0016]FIG. 6 is a pseudo-ternary diagram as a function of chromium,molybdenum and vanadium at 0.55 weight percent carbon with regard toalloys of the invention at 1000° C.; and

[0017]FIG. 7 is a graph comparing hardness of the steel alloys of theinvention with conventional carburized alloys.

[0018]FIG. 8 is a graph containing Falex wear test data for steel alloysof the invention in comparison with conventional 8620 steel.

[0019]FIG. 9 is a graph of NTN 3 ball-on-rod rolling contact fatiguedata for alloys of the invention in comparison with conventional M50bearing steel.

DESCRIPTION OF THE PREFERRED EMBODIMENTS

[0020] The steel alloys of the invention were developed using variousmodeling techniques followed by experimental confirmation or testing. Animportant component of the modeling is the application of athermochemical data bank and software system. The system or programemployed uses thermodynamic assessments from binary, ternary, andquaternary systems to extrapolate to higher order multicomponentsystems. Equilibria, constrained equilibria, and driving forces can becalculated as functions of composition, chemical potential, as well asother user defined functions. To apply this information to the modelingof highly nonequilibrium processes of interest in real alloys, thedynamic nature of phase transformations in terms of thermodynamicscaling factors are described and then evaluated by the Thermochemicalsoftware. Thus, hypothetical steel compositions were the subject of aninitial computational model involving the precipitation of M₂C carbidesleading to a secondary hardening response in ultrahigh-strength steels.A second effort employed a published thermodynamics-based model for thenon-linear composition dependence of the martensite start temperature. Athird modeling effort involves the application of quantum mechanicalcalculations to the production of hypothetical compositions with thegoal of achieving improved resistance to hydrogen embrittlement andintergranular fracture. Modeling techniques were then followed bytesting of the optimized alloys. Following is a discussion of modelingtechnique considerations.

[0021] Secondary Hardening

[0022] Ultrahigh-strength (UHS) secondary hardening steels arestrengthened by the precipitation of coherent M₂C carbides duringtempering. In high Co steels in which dislocation recovery is retarded,the M₂C carbides precipitate coherently on dislocations and provide thecharacteristic secondary hardening peak during tempering. A wide rangeof techniques are utilized to gather experimental information across acomplete range of size and time scales of interest. Atom-probe field-ionmicroscopy (APFIM), transmission electron microscopy (TEM), small angleneutron scattering (SANS), and X-ray diffraction (XRD) techniquesprovide information on the size, shape, composition, and latticeparameters of the M₂C precipitates as well as the resulting hardnessvalues spanning tempering times of less than an hour to more than athousand hours. This study identified that the precipitation was welldescribed by a theory developed by Langer and Schwartz for precipitationat high supersaturation in which the growth regime is suppressed andprecipitation occurs by a process of nucleation and coarsening,maintaining a particle size close to the critical size.

[0023] Based on these investigations, two important scaling factors areidentified. The initial critical nucleus determines the size scale ofthe precipitates throughout the precipitation reaction and thecoarsening rate constant determines the precipitation time scale. Thepeak hardness in an ultra high strength steel commonly occurs at theparticle size corresponding to the transition from particle shearing toOrowan bypass. It is also advantageous to bring the M₂C precipitation tocompletion in order to dissolve all of the transient cementite whichotherwise limits toughness and fatigue life. Therefore, the smaller theinitial critical particle size, the closer completion of precipitationoccurs to peak hardness and more efficient strengthening is obtained.The time scale of precipitation is also important due to the kineticcompetition between the secondary hardening reaction and the segregationof impurities to the prior austenite grain boundaries leading tointergranular embrittlement.

[0024] The initial critical nucleus size scales inversely with thethermodynamic driving force for precipitation. In the case of the M₂Ccarbide it is important to include the influence of prior cementiteformation and coherency on this quantity. The coherency elastic selfenergy can be evaluated by the calculation of an anisotropic ellipsoidalinclusion using the equivalent Eigenstrain method and the impact ofsolute redistribution on the resulting stress distribution is addressedby using open-system elastic constants. By relating the coherency strainto composition via the compositional dependence of the particle andmatrix lattice parameters, the composition dependence of the elasticself energy is determined in a form compatible with the thermodynamicsoftware. The linear elastic self energy calculation represents an upperlimit and a correction factor is used to fit the precipitationcomposition trajectories of a large set of experimental alloys.

[0025] The impact of prior cementite precipitation is accounted for bythe calculation of the coherent driving force in the presence of thecarbon potential due to para-equilibrium cementite. Thispara-equilibrium carbon potential is defined by an equilibrium betweenthe matrix and cementite in which the substitutional species are heldconstant and only the interstitial carbon is allowed to partition. Inthis approximation the cementite acts as a carbon source at constantchemical potential.

[0026]FIG. 1 represents the level of agreement of the strengtheningresponse of the model alloys with the above model. The model alloyscontain 16 wt % Co, 5 wt % Ni, and 0.24 wt % C with varying amounts ofthe carbide formers Cr, Mo, and, in a few cases, V. The nickel contentis chosen to eliminate austenite precipitation during tempering whichotherwise complicates the hardening response. In FIG. 1, the peakhardness during tempering at 510° C. is plotted against the drivingforce for precipitation of the coherent M₂C carbide in the presence ofpara-equilibrium cementite. The open circles represent alloys containingV. The relationship demonstrates the ability to predict peak hardnessvalues within approximately +/−25 VHN in this class of alloys.

[0027] The time scale of precipitation at high supersaturations,according to the Langer-Schwartz treatment, scales with the coarseningrate of the particle distribution. The modeling pursued in this workexpands upon the Lifshitz-Slyozov and Wager (LSW) theory, describing thecoarsening of spherical particles in a binary system, with the intent ofremoving the binary restrictions of the LSW theory and reformulating itin a manner compatible with the multicomponent thermodynamics of thesoftware and data bank system.

[0028] The result of this analysis characterizes the coarsening rate ofa particle of average size as a function of the multicomponent diffusioncoefficients, the equilibrium partitioning coefficients, and the secondderivatives of the Gibbs free-energy evaluated at the equilibrium state.The surface energy and molar volume are taken to be compositionindependent and are considered constant. In this form, the coarseningrate constant is the result of an asymptotic analysis and is onlyrepresentative at very long time scales and very close to equilibrium.This is certainly not the case for the precipitation of the M₂C carbideat high supersaturation. The matrix content of alloy is far fromequilibrium during much of the precipitation process, approachingequilibrium only near completion. This effect is more severe for alloyscontaining stoichiometric quantities of carbide formers as measured bythe relative difference in the matrix alloy content during precipitationand at equilibrium.

[0029] During precipitation in a stoichiometric alloy, the alloy matrixcontent is of the same order as the overall alloy content, while atequilibrium, the matrix alloy content is very small. To define acoarsening rate constant more representative of the conditions presentduring the precipitation process, a coarsening rate is evaluated at thepoint when the volume fraction of the precipitate is one-half of theequilibrium value. This is achieved by calculating the coherentequilibrium for the M₂C carbide, and then, adding energy to the M₂Cphase to account for capillarity until the amount of the phase is halfof the equilibrium value. The coarsening rate is then calculated fromthe thermodynamic properties of this state. FIG. 2 represents thecorrelation between the precipitation half-completion time and thehalf-completion coarsening rate constant of the model alloys for whichthis data is available.

[0030] Ms Temperatures

[0031] To predictively control the spatial distribution ofmartensite-start (Ms) temperatures in the carburized steels to achievefully martensitic structures with controlled residual stressdistributions, a published model was employed. The thermodynamics-basednucleation-kinetic model was calibrated to the composition-dependence ofmeasured Ms temperatures using both literature data and assessments ofexperimental multicomponent alloys.

[0032] Interfacial Cohesion

[0033] Intergranular embrittlement phenomena such as hydrogenembrittlement are undesirable in the intended alloys. Embrittlement ofultrahigh-strength steels is associated with the prior segregation tothe grain boundary of impurities such as P and S. A thermodynamictreatment of this phenomenon by Rice and Wang illustrates that thepotency of a segregating solute in reducing the work required forbrittle fracture along a boundary is linearly related to the differencein the segregation energy for the solute at the boundary and at the freesurface. Specifically, a solute with a higher segregation energy at thefree surface will be an embrittler while a solute with a highersegregation energy at the grain boundary will enhance intergranularcohesion. A survey of reported segregation energies and embrittlingpotency (reported as the shift in the ductile-to-brittle transitiontemperature per atomic percent solute on the grain boundary) in Fe-basealloys demonstrates these general trends; however, the experimentaldifficulty of surface thermodynamic measurements gives ambiguous valuesfor some solutes.

[0034] First principle calculations were used to determine the totalenergy of atomic cells representing the Fe Σ3 [110](111) grain boundaryand (111) free surface with a monolayer of an impurity solute present.The calculations were accomplished with the full-potential linearizedplane wave (FLAPW) total energy technique. The atomic structure in eachcase was relaxed to find the minimum energy state. The results of thesecalculations include not only the segregation energies responsible forthe embrittling or cohesion enhancing effects of segregating solutes,but the underlying electronic structure of the solutes in the boundaryand surface environments. A comparison of the directional covalentelectronic structure between B, a strong cohesion enhancer, and P, astrong embrittler, indicates the strong bonding of the B atom across theboundary plane associated with hybridization of the B 2p electrons withthe Fe d band. This directional bonding is not seen in the case of the Patom which does not significantly hybridize with Fe.

[0035] The results of the first principle calculations were correlatedto the experimental embrittling potency in FIG. 3. The differencebetween grain boundary and free surface segregation energies, calculatedby electronic structure calculations, and the experimentally observedshift in the ductile-to-brittle-transition-temperature are plotted forC, B, P and S solutes. The C and B are shown as cohesion enhancers, Pand S as embrittlers. The computed energy differences are in excellentagreement with the observed effects on interfacial cohesion.

[0036] Materials Design

[0037] Background

[0038] Design considerations for high performance gears in aerospace,automotive, and other applications include the desire to transmit morepower in less space and with less gear weight. Current high performancegear steels are typically stage I quench and tempered martensites, casehardened through carburizing to 60 R_(c) on the surface with a corehardness of typically 35 R_(c) as gear teeth are typically subjected tobending and contact stresses. Failure modes in gears are generallygrouped into three categories, tooth breakage, surface pitting, andsubsurface spalling. High surface hardness is used to limit the toothbreakage due to bending fatigue as well as pitting failures. If thecontact fatigue strength drops below the applied stress at any pointbelow the surface, subsurface spalling can result. Typically a 1 mm casedepth is required to provide adequate fatigue strength to avoidsubsurface spalling.

[0039] Stage IV tempered secondary hardening steels offer numerousadvantages over conventional gear steels. The efficient strengtheningand superior toughness achieved in secondary hardening steels allowsgreater core and case hardness to be attained reducing the size andweight of gears needed to transmit a given power. The greatertemperature resistance of secondary hardening steels allows operation athigher temperatures and longer times before performance degrades throughmechanisms such as scoring. In addition, recent results on commercialsecondary hardening alloys indicate casting of these steels may bepossible.

[0040] Analysis

[0041] The systems analysis of the case-core secondary hardening steelsystem is the first step in the design process. FIG. 4 illustrates thetotal processing/structure/properties/performance system structure forhigh power-density gears manufactured by three alternative processingroutes, conventional forged ingot processing, near net shape casting andpowder processing. Case hardenable secondary hardening gear steels are asubsystem of this flow-block diagram and are the focus of thisdisclosure. The sequential processing steps dictate the evolution of thecase and core microstructures which determine the combination ofproperties required for the overall performance of the system. Both thecase and core consist of microstructures of a martensite with high Co,for dislocation recovery resistance essential for efficient secondaryhardening, and Ni, for cleavage resistance. Strengthening is provided bythe coherent precipitation of fine M₂C carbides on dislocations. Thissecondary hardening reaction dissolves the transient cementite and itbrings the precipitation reaction to completion in order to eliminatecementite for high toughness and fatigue resistance. The grain refiningdispersion has a double impact on toughness. By limiting grain growth athigh temperature during solution treatment, brittle intergranular modesof fracture are inhibited.

[0042] The grain refining particles also play an important role in theductile microvoid nucleation and coalescence fracture behavior. Thus, itis desired to preserve adequate volume fraction and size to pin thegrain boundaries while choosing the phases with greatest interfacialcohesion. Also desirable is the control of the grain boundary chemistryto avoid intergranular embrittlement (such as by hydrogen embrittlement)in association with prior segregation of embrittling impurities. Duringtempering, impurities segregate to the grain boundaries and in the caseof P and S reduce the interfacial cohesion of the boundary promotingintergranular embrittlement. A number of methods are used to avoid thisproblem. Gettering compounds can be utilized to tie up the impurities instable compounds reducing the segregation to the grain boundary. Inorder to produce the most stable compounds, however, rapidsolidification processing is required. Within their solubility limits,additional segregants such as B can be deliberately added to enhanceintergranular cohesion, and the precipitation rate for the secondaryhardening reaction can be increased to limit the time at temperature forharmful grain boundary segregation.

[0043] Design

[0044] As a first design step, core and case carbon levels required forthe desired hardness are estimated. This is done by fitting data forexisting secondary hardening Ni—Co steels to an Orowan strengtheningmodel and extrapolating to the desired strength. It is estimated that acore carbon content of 0.25 wt % and a case carbon content of 0.55 wt %is needed to provide the desired core and case hardness in this Ni—Costeel.

[0045] The next step is to determine the matrix composition of the Fe,Ni, and Co. In order to produce the desired lath martensite morphology,an M_(s) temperature of 200° C. or above is required. Using thenucleation kinetic model for the compositional dependence of the M_(s)temperature, the variation with Ni and Co content is determined. Thisresult is illustrated in FIG. 5 for the case carbon content using apreliminary composition of carbon formers equal to 5 wt % Cr, 0.5 wt %Mo and 0.0 wt % V. Since the case has a higher carbon content than thecore, the core will possess a higher M_(s) temperature than the case. InFIG. 5 the Co and Ni content required to fix the M_(s) temperature at200° C. is indicated. Since a high Ni content is desired to avoidcleavage fracture, a composition containing 25 wt % Co was chosen. Thisallows the highest possible Ni content, approximately 3.5 wt %, to beused. These calculations are later repeated for consistency when thecomposition of the carbide formers is further refined.

[0046] To define the optimal composition of the carbide formers a numberof design constraints are applied. The total amount of carbide formersin the alloy must be greater than that required to consume the carbonpresent in the case. This lower limit insures that, at completion,embrittling cementite is completely converted to M₂C carbide. In orderto reduce grain boundary segregation, the precipitation rate ismaximized. This allows the shortest possible tempering time. Thecoherent precipitation driving force is maximized to provide a smallcritical particle size for the M₂C and more efficient strengthening.Finally, the solution temperature is limited to 1000° C. This allows Cr,Mo and V containing carbides such as M₂₃C₆, M₇C₃, MC and M₆C to bedissolved at reasonable processing temperatures while maintaining veryfine scale TiC carbides to act as the grain refining dispersion.

[0047] Calculations for the precipitation rate constant indicate low Mocompositions are favorable, while driving force calculations havedemonstrated the highly beneficial effect of higher V contents. Thesolubility constraints are presented by the diagram in FIG. 6. Here theequilibrium phase fields at 1000° C. are given as a function of Cr and Vcontent. The Mo content is determined by the stoichiometry requirements,the matrix composition is taken from the earlier calculations, and thecarbon content represents the case composition. The point on the diagramwithin the single phase FCC field that maximizes the V content andminimizes the Mo content represents the compromise fulfilling the designcriteria. This composition is 4.8 wt % Cr, 0.03 wt % Mo, and 0.06 wt %V. A recalculation of the matrix composition using the final carbideformers results in an alloy composition of Fe-25 Co-3.8 Ni-4.8 Cr-0.03Mo-0.06 V-0.55 (case)/0.25 (core) C (in wt %). Consistent with the modelpredictions of FIG. 3, a soluble boron addition of 15-20 ppm is added toenhance intergranular embrittlement resistance.

EXAMPLES

[0048] A 17 lb. vacuum induction heat of the above composition wasprepared from high purity materials. The ingot was forged at 1150° C. ina bar 1.25″ square by 38″ long. The M_(s) temperature of the alloy wasdetermined from dilatometery and found to agree with model predictions.The solution treatment response of the alloy was determined fromhardness measurements in the stage I tempered condition. The optimumprocessing conditions for the core material was determined to be a 1050°C. 1 hour solution treatment followed by an oil quench and a liquidnitrogen deep freeze. After optimal solution treatment, a 12 hour temperat 482° C. results in the desired overaged hardness of 55 R_(c) for thecore material. The material was then plasma carburized and processedusing these parameters. The C potential, temperature and time used inthe carburizing treatment were determined from simulations withmulticomponent diffusion software to provide the target surface carboncontent of 0.55 wt % and a 1 mm case depth . The curve labeled C2 inFIG. 7 represents the hardness profile achieved for the carburizedsample. A surface hardness of 67 HR_(c) and a case depth of 1 mm areobtained.

[0049] Using techniques and processes of this nature, the followingalloys set forth in Table 1 were developed and tested: TABLE 1 Alloy FeCo Ni Cr Mo V C (Core) A1 Bal. 18 9.5 3.5 1.1 0.08 0.20 C2 Bal. 25 3.84.8 0.03 0.06 0.237 C3 Bal. 28 3.25-3.15 5.0 1.75-2.50 0.025 0.05-0.18CS1 Bal. 15 1.5 9.0 0.0 0.2 0.05-0.20

[0050] The first, A1, is targeted as a replacement for current gearmaterials in applications where component redesign is not feasible buthigher core strength and toughness is needed. As such, A1 has surfacewear properties similar to current commercial properties, but possessessuperior core toughness and strength 54 HRC and a K_(IC) of >75Ksi{square root}in. The second alloy C2 corresponds to the prototypealloy just described. The third alloy, C3, pushes the surface propertiess far as possible while maintaining adequate core strength andtoughness. As also shown by the hardness profiles of FIG. 7, the alloyhas reached a surface hardness corresponding to HRC 69. Wear tests forthe carburized material in a standard Falex gear simulator show muchreduced weight loss compared to standard carburized 8620 steel in FIG.8. A ball-on-rod rolling contact fatigue test (NTN type) conducted at786 ksi Hertzian contact stress indicates an order of magnitude increasein L₁₀ fatigue life compared to M50 bearing steel as shown in FIG. 9.The fourth alloy in Table 1, CS1, represents a stainless variant of thisclass of alloy. Targeted to match the surface properties of standardnon-stainless gear and bearing materials with sufficient core strengthand toughness, the alloy has achieved corrosion resistance better than440C by anodic polarization conducted in distilled water with a neutralph. (sucrose added for electrical conductivity). Similar relativebehaviors were demonstrated in 3.5% NaCl solution. In salt fog tests,CS1 outperformed 440C and commercial carburizing stainless steels, theperformance gap widened when the tests were completed on samples in thecarburized condition. The carburized alloy achieved surface mechanicalproperties equivalent to A1 while maintaining corrosion resistance. InRFC tests of the type represented in FIG. 9, both A1 and CS1 showed L₁₀fatigue life equal or superior to the M50 bearing steel. The table(Table 2) below summarizes the performance achieved in the four alloys:TABLE 2 Rolling Contact Core Core Surface Bending Fatigue⁵ CorrosionAlloy Hardness¹ Toughness² Hardness³ Fatigue⁴ L₁₀ Resistance⁶ A1  54R_(C) >75 ksi{square root}in >61 R_(C) ≧EN36C ≧M50 NA C2 ≦58 R_(C)adjustable  67 R_(C) NA in process NA C3 ≦59 R_(C) adjustable  69 R_(C)NA 10 × M50 NA CS1 ≦53 R_(C) ≧25 ksi{square root}in  63 R_(C) NA≧M50 >440C adjustable

[0051] Each alloy has a surface hardness exceeding and core hardnessexceeding prior art compositions achieved at a lower cost. The alloys ofthe invention are compounded and made in accord with this disclosure,but variants within the skill of the art are possible. The invention istherefore limited only by the following claims and equivalents.

What is claimed is:
 1. A steel alloy comprising, in combination: a)about 0.05 to about 0.24 weight percent carbon; b) about 15.0 to about28.0 weight percent cobalt; c) about 1.5 to about 9.5 weight percentnickel; d) one or more additives selected from a group consisting of: 1)about 3.5 to about 9.0 weight percent chromium; 2) up to about 2.5weight percent molybdenum; 3) up to 0.2 weight percent vanadium; 4) andthe balance, iron. said alloy having a case and a core, said case havinga carbon content greater than said core, said alloy comprising amartensite phase and having a case hardness of at least about Rockwell60 C.